WO2011111330A1 - High-strength steel sheet and method for producing same - Google Patents

Abstract

Through the disclosed method, it is possible to obtain a highly ductile and high tensile strength high-strength steel sheet having a tensile strength of at least 1470 MPa and a tensile strength × total elongation of at least 29000 MPa · % by means of: a steel composition containing at least 0.30% and no more than 0.73% C, no more than 3.0% Si, no more than 3.0% Al, at least 0.7% Si+Al, at least 0.2% and no more than 8.0% Cr, no more than 10.0% Mn, at least 1.0% Cr+Mn, no more than 0.1% P, no more than 0.07% S, and no more than 0.010% N, the remainder being Fe and unavoidable impurities; and a steel structure of which the area ratio of martensite to the total steel sheet structure is 15-90% inclusive, the amount of residual austenite is 10-50% inclusive, at least 50% of said martensite is tempered martensite, the area ratio of said tempered martensite to the total steel sheet structure is at least 10%, and the area ratio of polygonal ferrite to the total steel sheet structure is 10% or less (including 0%).

Description

High strength steel plate and manufacturing method thereof

The present invention relates to a high-strength steel sheet having a tensile strength (TS) of 1470 MPa or more excellent in workability, particularly ductility, and a method for producing the same used in industrial fields such as automobiles and electrical equipment.

In recent years, improving the fuel efficiency of automobiles has become an important issue from the viewpoint of global environmental conservation. For this reason, efforts are being made to reduce the thickness of vehicle body parts by increasing the strength of vehicle body materials and to reduce the weight of the vehicle body itself.

Generally, in order to increase the strength of a steel sheet, it is necessary to increase the ratio of hard phases such as martensite and bainite to the entire structure of the steel sheet. However, since increasing the strength of the steel sheet by increasing the proportion of the hard phase causes a decrease in workability, development of a steel sheet having both high strength and excellent workability is desired. So far, various composite steel sheets have been developed, such as ferritic-martensitic duplex steel (DP steel) and TRIP steel using transformation-induced plasticity of retained austenite.

When increasing the ratio of the hard phase in the composite structure steel plate, the workability of the steel plate is strongly influenced by the workability of the hard phase. This is because when the ratio of hard phase is small and soft polygonal ferrite is large, the deformability of polygonal ferrite dominates the workability of the steel sheet, and even when the hard phase has insufficient workability. While workability such as ductility is secured, when the ratio of the hard phase is large, the deformability of the hard phase itself directly affects the formability of the steel sheet, not the deformability of polygonal ferrite. It is.

For this reason, in the case of a cold-rolled steel sheet, after performing annealing and adjusting the amount of polygonal ferrite generated in the subsequent cooling process, the steel sheet is water-quenched to generate martensite, and the steel sheet is raised again. By heating and holding at a high temperature, the martensite has been tempered to generate carbides in the martensite that is a hard phase, thereby improving the workability of the martensite. However, normally, in the case of continuous annealing water quenching equipment that performs such water quenching, since the temperature after quenching is necessarily close to the water temperature, most of the untransformed austenite undergoes martensitic transformation. And other low-temperature transformation structures were difficult to use. Therefore, the improvement of the workability of the hard structure is limited to the effect of tempering the martensite, and as a result, the improvement of the workability of the steel sheet is also limited.

Further, as steel plates having a hard phase other than martensite, there are steel plates in which the main phase is polygonal ferrite, the hard phase is bainite or pearlite, and carbides are generated in these hard phases bainite or pearlite. This steel sheet is a steel sheet that not only improves the workability with polygonal ferrite alone, but also improves the workability of the hard phase itself by generating carbides in the hard phase, and in particular, improves the stretch flangeability. . However, since the main phase is polygonal ferrite, it has been difficult to achieve both high strength exceeding 1180 MPa in tensile strength (TS) and workability.

With respect to a composite structure steel sheet containing retained austenite, for example, Patent Document 1 specifies an alloy component and makes the steel structure fine and uniform bainite having retained austenite, which is excellent in bending workability and impact properties. Tensile steel sheets have been proposed.
Patent Document 2 discloses a composite steel sheet having excellent bake hardenability by defining predetermined alloy components, making the steel structure bainite having retained austenite, and defining the amount of retained austenite in bainite. Proposed.
Further, Patent Document 3 defines predetermined alloy components, the steel structure is 90% or more in area ratio of bainite having retained austenite, the amount of retained austenite in bainite is 1% or more and 15% or less, and bainite. By defining the hardness (HV) of the steel sheet, a composite structure steel plate excellent in impact resistance has been proposed.

JP-A-4-235253 JP 2004-76114 A Japanese Patent Laid-Open No. 11-256273

However, the above-described steel sheet has the following problems.
In the component composition described in Patent Document 1, it is difficult to secure a stable amount of retained austenite that exhibits a TRIP effect in a high strain region when strain is applied to a steel sheet, and bendability is obtained. However, the ductility until plastic instability occurs is low, and the stretchability is inferior.
Although the steel sheet described in Patent Document 2 has bake hardenability, it has a structure in which bainite or ferrite is mainly contained and martensite is suppressed as much as possible, so that it has a tensile strength (TS) exceeding 1180 MPa. It is also difficult to ensure workability at the time of increasing strength.
The main purpose of the steel sheet described in Patent Document 3 is to improve impact resistance. Since the main phase is bainite having a hardness of HV250 or less, specifically, it is a structure containing more than 90%, It is extremely difficult to make the tensile strength (TS) exceed 1180 MPa.

On the other hand, among automotive parts molded by pressing, for example, steel plates used as materials for parts that require particularly high strength, such as door impact beams and bumper reinforcements that suppress deformation in the event of a car crash, and more in the future Furthermore, it is considered that a tensile strength (TS) of 1470 MPa or more is required. In addition, structural parts such as members and center pillar inners, which are structural parts with relatively complicated shapes, are expected to have a tensile strength (TS) of 980 MPa or higher, and more than 1180 MPa in the future.

The present invention advantageously solves the problem that it has been difficult to ensure workability because of its high strength so far, and it is advantageous to produce a high-strength steel sheet with a tensile strength (TS) of 1470 MPa or more and excellent ductility. It is intended to be provided with a method.
Moreover, the high-strength steel sheet of the present invention includes a steel sheet obtained by subjecting the surface of the steel sheet to hot dip galvanization or galvannealing.
In the present invention, excellent workability means that the value of tensile strength (TS) × total elongation (T.EL) is 29000 MPa ·% or more.

The inventors have made extensive studies on the component composition and microstructure of the steel sheet in order to solve the above problems. As a result, the martensite structure is utilized to increase the strength, the C content in the steel sheet is increased to 0.30% or more and the C content is increased, and the ferrite formation suppressing effect and the workability improving effect during martensite tempering are also achieved. After adding Cr having, after quenching the steel sheet annealed in the austenite single phase region and partially martensite transformation of the austenite, by tempering martensite and stabilizing the retained austenite, workability, especially It was found that a high-strength steel sheet having a remarkably excellent balance between strength and ductility and having a tensile strength of 1470 MPa or more can be obtained.

The present invention is based on the above findings, and the gist of the present invention is as follows.
(1) By mass% C: 0.30% or more and 0.73% or less,
Si: 3.0% or less,
Al: 3.0% or less,
Si + Al: 0.7% or more,
Cr: 0.2% to 8.0%,
Mn: 10.0% or less,
Cr + Mn: 1.0% or more
P: 0.1% or less,
S: 0.07% or less and N: 0.010% or less, the balance is composed of Fe and inevitable impurities,
As the steel sheet structure, the area ratio of martensite to the entire steel sheet structure is 15% to 90%, the amount of retained austenite is 10% to 50%, and 50% or more of the martensite is tempered martensite and the tempered martensite. Satisfies the area ratio of the site steel sheet structure to 10% or more, the area ratio of polygonal ferrite to the entire steel sheet structure of 10% or less (including 0%), tensile strength of 1470 MPa or more, tensile strength x total elongation Is a high-strength steel sheet characterized by a 29000 MPa ·% or more.

(2) 30% or more of the total length of the prior austenite grain boundary is present in the tempered martensite or is adjacent to the tempered martensite. Strength steel plate.

(3) The high-strength steel sheet according to (1) or (2) above, wherein an average amount of C in the retained austenite is 0.7% by mass or more.

(4) The steel sheet is further in mass%,
Any one of the above (1) to (3), characterized in that it contains Ni: 0.05% or more and 5.0% or less and satisfies Cr + Mn + Ni: 1.0% or more instead of Cr + Mn: 1.0% or more High strength steel plate.

(5) The steel sheet is further mass%,
V: 0.005% to 1.0%,
Any one of the above (1) to (4), characterized by containing one or more selected from Mo: 0.005% to 0.5% and Cu: 0.05% to 2.0% The high-strength steel sheet described in 1.

(6) The steel sheet is further mass%,
Any one of (1) to (5) above, characterized by containing one or two selected from Ti: 0.01% to 0.1% and Nb: 0.01% to 0.1% High strength steel sheet as described.

(7) The steel sheet is further mass%,
B: The high-strength steel sheet according to any one of (1) to (6) above, which contains 0.0003% or more and 0.0050% or less.

(8) The steel sheet is further in mass%,
Any one of (1) to (7) above, characterized by containing one or two selected from Ca: 0.001% to 0.005% and REM: 0.001% to 0.005% High strength steel sheet as described.

(9) A high-strength steel sheet comprising a hot-dip galvanized layer or an alloyed hot-dip galvanized layer on the surface of the steel sheet described in any one of (1) to (8) above.

(10) A steel slab having the composition described in any one of (1) to (8) above is hot-rolled and then cold-rolled into a cold-rolled steel sheet, and then the cold-rolled steel sheet is austenite single-piece. After annealing in the phase range for 15 seconds or more and 1000 seconds or less, the sample is cooled at an average cooling rate of 3 ° C./s or more to the first temperature range of Ms−150 ° C. or more and less than Ms with respect to the martensite transformation start temperature Ms. A method for producing a high-strength steel sheet, characterized in that the temperature is raised to a second temperature range of 340 ° C. or more and 520 ° C. or less, and subsequently maintained in the second temperature range for 15 seconds or more and 1000 seconds or less.

(11) The hot-dip galvanizing treatment or the alloying hot-dip galvanizing treatment is performed during the temperature rise to the second temperature range or during the holding in the second temperature range. Manufacturing method of high strength steel sheet.

According to the present invention, it is possible to stably obtain a high-strength steel sheet having remarkably excellent workability, particularly ductility, and a tensile strength (TS) of 1470 MPa or more.
Therefore, the present invention has a very high utility value in the industrial field such as automobiles and electrical equipment, and is extremely useful particularly for reducing the weight of automobile bodies.

It is the figure which showed the temperature pattern of the heat processing in the manufacturing method according to this invention.

Hereinafter, the present invention will be specifically described.
First, the reason why the steel sheet structure is limited as described above in the present invention will be described. Hereinafter, the area ratio is the area ratio relative to the entire steel sheet structure.

Martensite area ratio: 15% or more and 90% or less Martensite is a hard phase and a structure necessary for increasing the strength of a steel sheet. When the area ratio of martensite is less than 15%, the tensile strength (TS) of the steel sheet does not satisfy 1470 MPa. On the other hand, if the area ratio of martensite exceeds 90%, a stable retained austenite amount cannot be secured, so that there is a problem that workability such as ductility is lowered. Therefore, the area ratio of martensite is 15% or more and 90% or less. Preferably they are 20% or more and 80% or less.

Of martensite, tempered martensite ratio: 50% or more Tempered martensite area ratio: 10% or more The tempered martensite ratio is less than 50% of the total martensite area ratio or 10% of the entire steel structure. If it is less than 1, the tensile strength is 1470 MPa or more, but sufficient ductility may not be obtained. This is because as-quenched martensite containing high C is extremely hard, has low deformability and is poor in toughness, and when the amount increases, it becomes brittle when strain is applied and consequently no excellent ductility can be obtained. It is. By tempering such as-quenched martensite, although the strength is slightly reduced, the deformability of martensite itself is greatly improved, so brittle fracture does not occur at the time of applying strain, and the structure of the present invention By realizing the configuration, TS x T.EL can be over 29000MPa%. Therefore, the ratio of tempered martensite in martensite is 50% or more of the total martensite area ratio present in the steel sheet, and the area ratio of tempered martensite to the entire steel sheet structure is 10% or more. Preferably, 70% or more of the total martensite area ratio and 20% or more of the area ratio relative to the entire steel sheet structure, more preferably 80% or more of the total martensite area ratio and 30% of the total area ratio of the steel sheet That's it. Tempered martensite is observed as a microstructure in which fine carbides are precipitated in martensite by observation with a scanning electron microscope (SEM), and as-quenched martensite in which such carbides are not observed inside martensite. It can be clearly distinguished from the site.

Residual austenite amount: 10% or more and 50% or less Residual austenite is transformed into martensite by the TRIP effect during processing, and the strength is increased by hard martensite containing high C and at the same time the ductility is improved by increasing the strain dispersibility. Improve.
In the steel sheet of the present invention, after partially martensitic transformation is performed, for example, the upper austenite transformation that suppresses the formation of carbides is used to form retained austenite having a particularly high carbon concentration. As a result, retained austenite that can exhibit the TRIP effect even in a high strain region during processing can be obtained. In the present invention, it is important to secure a predetermined amount of stable retained austenite having a high carbon concentration, and the effective utilization of the upper bainite transformation is to suppress the formation of carbides. However, utilization of this upper bainite transformation is not necessarily essential. For example, in a state where the martensite fraction is high, it is possible to concentrate carbon in austenite while maintaining a high temperature after quenching.
By utilizing such retained austenite and martensite together, good workability can be obtained even in a high strength region where the tensile strength (TS) is 1470 MPa or more. Specifically, TS × T.EL Can be made 29000 MPa ·% or more, and a steel sheet with an extremely excellent balance between strength and ductility can be obtained.

Here, since retained austenite is distributed in a state surrounded by tempered martensite, it is difficult to accurately quantify the amount (area ratio) by microstructure observation, but the amount of retained austenite conventionally used is measured. The intensity measurement method by X-ray diffraction (ERD), which is a technique to be used. Specifically, if the retained austenite amount obtained from the X-ray diffraction intensity ratio of ferrite and austenite is 10% or more, a sufficient TRIP effect can be obtained, and the tensile strength (TS) is 1470 MPa or more. It has been confirmed that T.EL can achieve 29000 MPa ·% or more. It has been confirmed that the amount of retained austenite obtained by a conventional method for measuring the amount of retained austenite is equivalent to the area ratio of retained austenite to the entire steel sheet structure.

If the amount of retained austenite is less than 10%, sufficient TRIP effect cannot be obtained. On the other hand, if it exceeds 50%, the hard martensite generated after the TRIP effect is manifested becomes excessive, and deterioration of toughness becomes a problem. Therefore, the amount of retained austenite is in the range of 10% to 50%. Preferably, it is in the range of 14% to 45%. More preferably, it is in the range of 18% or more and 40% or less.

Polygonal ferrite area ratio: 10% or less (including 0%)
If the area ratio of polygonal ferrite exceeds 10%, it will be difficult to satisfy the tensile strength (TS) of 1470 MPa or more, and at the same time, strain will concentrate on the soft polygonal ferrite mixed in the hard structure during processing. Therefore, cracks are easily generated during processing, and as a result, desired workability cannot be obtained. Here, if the area ratio of polygonal ferrite is 10% or less, even if polygonal ferrite is present, a small amount of polygonal ferrite is isolated and dispersed in the hard phase, and strain concentration can be suppressed. Degradation of workability can be avoided. Therefore, the area ratio of polygonal ferrite is 10% or less. Preferably it is 5% or less, More preferably, it is 3% or less, and 0% may be sufficient.

The steel sheet of the present invention may contain pearlite, Widmanstatten ferrite, and lower bainite as the remaining structure as well as upper bainite that may be produced after part of martensite. In that case, the allowable content of the remaining structure excluding the upper bainite is preferably 20% or less in terms of area ratio. More preferably, it is 10% or less. On the other hand, upper bainite is a structure that may occur when tempering martensite as it is quenched, and when its content is excessive, it is difficult to ensure strength particularly exceeding 1.7 GPa, so the area ratio relative to the entire structure 60% or less is preferable. More preferably, it is less than 50%, More preferably, it is less than 35%.

The above is the basic structure of the steel sheet structure in the high-strength steel sheet of the present invention, but the following structure may be added as necessary.
More than 30% of the total length of the prior austenite grain boundary is present in tempered martensite or adjacent to tempered martensite. When the steel has a structure composed of high austenite retained martensite or martensite as in the present invention steel, the steel itself is high. Due to its strength, fracture may occur from the prior austenite grain boundaries during molding and processing. This is thought to be caused by the lack of toughness of the prior austenite grain boundaries, but by allowing the prior austenite grain boundaries to exist inside tempered martensite with excellent workability or adjacent to tempered martensite, Molding / workability can be improved.
In order to obtain such an effect, 30% or more of the total length of the prior austenite grain boundary needs to be present in the tempered martensite or adjacent to the tempered martensite. Preferably it is 45% or more. The total length of prior austenite grain boundaries can be determined from the length of prior austenite grain boundaries revealed by the technique disclosed in Japanese Patent Application Laid-Open No. 2005-241635. Further, the region of the same visual field is buffed again and subjected to nital corrosion, whereby the ratio of the former austenite grain boundary existing in the tempered martensite or adjacent to the tempered martensite can be obtained.

Average C content in retained austenite: 0.70% or more To obtain excellent workability by utilizing the TRIP effect, C content in retained austenite is required for high-strength steel sheets with a tensile strength (TS) of 1470 MPa or more. is important. As a result of investigations by the inventors, in the steel sheet of the present invention, X-ray diffraction (XRD), which is a conventional method for measuring the average C content in retained austenite (average of C content in retained austenite), is performed. It has been found that if the average C content in the retained austenite obtained from the shift amount of the diffraction peak is 0.70% or more, further excellent workability can be obtained. If the average C content in the retained austenite is less than 0.70%, martensitic transformation may occur in the low strain region during processing, and the TRIP effect in the high strain region that improves workability may not be obtained sufficiently. . Therefore, the average C content in the retained austenite is preferably 0.70% or more, and more preferably 0.90% or more. On the other hand, if the average amount of C in the retained austenite exceeds 2.00%, the retained austenite becomes excessively stable, the martensitic transformation does not occur during processing, and the TRIP effect is not manifested, so there is a concern that the ductility may be lowered. Accordingly, the average C content in the retained austenite is preferably 2.00% or less.

Next, the reason why the component composition of the steel sheet is limited as described above in the present invention will be described.
In addition,% showing the following component compositions shall mean the mass%.
C: 0.30% or more and 0.73% or less C is an element indispensable for increasing the strength of a steel sheet and securing a stable retained austenite amount, and is necessary for securing a martensite amount and retaining austenite at room temperature. It is an element. If the C content is less than 0.30%, it is difficult to ensure the strength and workability of the steel sheet. On the other hand, if the amount of C exceeds 0.73%, the welded portion and the weld heat affected zone are hardened and the weldability deteriorates. Therefore, the C content is in the range of 0.30% to 0.73%. Preferably, it is in the range of more than 0.34% and 0.69% or less, and more preferably 0.39% or more.

Si: 3.0% or less Si is a useful element that contributes to improving the strength of steel by solid solution strengthening. However, if the amount of Si exceeds 3.0%, the workability and toughness will deteriorate due to the increase in the solid solution amount in polygonal ferrite, and surface properties will deteriorate due to the occurrence of red scale, etc. May cause deterioration of plating adhesion and adhesion, so the Si content is 3.0% or less. More preferably, it is 2.6% or less, more preferably 2.2% or less, and may be 0%.

Al: 3.0% or less Al is a useful element that is added as a deoxidizer in the steelmaking process. However, if it exceeds 3.0%, inclusions in the steel sheet increase and ductility may be deteriorated. Is 3.0% or less. More preferably, it is 2.0% or less. On the other hand, in order to obtain the deoxidation effect of Al, the Al content is preferably 0.001% or more, more preferably 0.005% or more. The amount of Al in the present invention is the amount of Al contained in the steel sheet after deoxidation. When deoxidizing with Si or the like, Al may be 0%.

Si + Al: 0.7% or more Both Si and Al are elements useful for suppressing the formation of carbides and promoting the generation of retained austenite, which is an important structure in securing the balance between strength and ductility in the present invention. . Although suppression of carbides is effective even if Si or Al is contained alone, it is necessary to contain at least 0.7% in total of the Si amount and Al amount.

Cr: 0.2% or more and 8.0% or less Cr is an essential element in the present invention, and has the effect of suppressing the formation of ferrite and pearlite during cooling from the annealing temperature, and at the same time improves the workability of martensite. Although the mechanism is not clear, it is thought that by changing the formation state of carbide, etc., even hard and high-strength martensite is realized to have excellent workability, and the effect is Cr The amount is obtained at 0.2% or more. Preferably it is 0.5% or more, More preferably, it is 1.0% or more. On the other hand, if the Cr content exceeds 8.0%, the amount of hard martensite becomes excessive, and the strength may be higher than necessary or sufficient ductility may not be obtained. Therefore, the Cr content is 8.0% or less. Preferably it is 6.0% or less, More preferably, it is 4.0% or less.

Mn: 10.0% or less Mn is an element effective for strengthening steel, and is preferably contained at 0.01% or more, and can be used together with Cr. However, if the content exceeds 10.0%, the castability deteriorates. Therefore, the amount of Mn needs to be 10.0% or less. Preferably it is 7.0% or less, More preferably, it is 4.0% or less. Note that Mn may be 0% when Cr or the like is fully utilized.

Cr + Mn: 1.0% or more Cr and Mn are elements that suppress the formation of ferrite, pearlite, and bainite during cooling from the annealing temperature. In the present invention, it is desirable to partially martensite the austenite generated during annealing while maintaining it as much as possible, and in order to realize this, the Cr + Mn content needs to be 1.0% or more. Preferably it is 1.5% or more.

P: 0.1% or less P is an element useful for strengthening steel. However, when the P content exceeds 0.1%, impact resistance deteriorates due to embrittlement due to grain boundary segregation, and alloyed molten zinc is added to the steel sheet. In the case of plating, the alloying speed is greatly delayed. Therefore, the P content is 0.1% or less. Preferably it is 0.05% or less. The amount of P is preferably reduced, but if it is less than 0.005%, it causes a significant increase in cost, so the lower limit is preferably about 0.005%.

S: 0.07% or less Since S becomes an inclusion such as MnS and causes deterioration in impact resistance and cracking along the metal flow of the welded portion, it is preferable to reduce the amount of S as much as possible. However, excessively reducing the amount of S causes an increase in manufacturing cost, so the amount of S is set to 0.07% or less. Preferably it is 0.05% or less, More preferably, it is 0.01% or less. Note that, if S is less than 0.0005%, there is a large increase in manufacturing cost, so the lower limit is about 0.0005% from the viewpoint of manufacturing cost.

N: 0.010% or less N is an element that most deteriorates the aging resistance of steel, and is preferably reduced as much as possible. When the N content exceeds 0.010%, deterioration of aging resistance becomes remarkable, so the N content is set to 0.010% or less. Note that, if N is less than 0.001%, a large increase in manufacturing cost is caused, so that the lower limit is about 0.001% from the viewpoint of manufacturing cost.

Moreover, in this invention, the component described below other than the above-mentioned basic component can be contained appropriately.
Ni: 0.05% or more and 5.0% or less, and Cr + Mn: 1.0% or more instead of Cr + Mn + Ni: 1.0% or more Ni, like Cr and Mn, is an element that suppresses the formation of ferrite, pearlite, and bainite when cooled from the annealing temperature. In order to obtain such an effect, the Ni content is preferably 0.05% or more. Further, in the present invention, as described above, it is desirable to partially martensite the austenite generated during annealing while maintaining it as much as possible. For that purpose, when Ni is contained, the amount of Ni is 0.05% or more. In addition, the Cr + Mn + Ni amount is preferably 1.0% or more instead of the above-described condition that the Cr + Mn amount is 1.0% or more. More preferably, the Ni amount is 0.05% or more and the Cr + Mn + Ni amount is 1.5% or more. Note that if the Ni content exceeds 5.0%, the workability of the steel sheet may be lowered, so the Ni content is preferably 5.0% or less.

V: 0.005% or more and 1.0% or less, Mo: 0.005% or more and 0.5% or less, Cu: 0.05% or more and 2.0% or less selected from among 0.05% and 2.0% or less V, Mo, and Cu are cooled at the annealing temperature. It is an element that has the effect of suppressing the formation of pearlite. The effect is obtained when V: 0.005% or more, Mo: 0.005% or more, and Cu: 0.05% or more. On the other hand, if it exceeds V: 1.0%, Mo: 0.5%, and Cu: 2.0%, the amount of hard martensite becomes excessive and the strength becomes higher than necessary. Therefore, when V, Mo, and Cu are contained, V: 0.005% to 1.0%, Mo: 0.005% to 0.5%, and Cu: 0.05% to 2.0%.

One or two types selected from Ti: 0.01% or more and 0.1% or less, Nb: 0.01% or more and 0.1% or less Ti and Nb are useful for the precipitation strengthening of steel, and the effect of each is 0.01% The above is obtained. On the other hand, if the respective contents exceed 0.1%, the workability and the shape freezing property decrease. Therefore, when Ti and Nb are contained, the range is Ti: 0.01% to 0.1% and Nb: 0.01% to 0.1%.

B: 0.0003% or more and 0.0050% or less B is an element useful for suppressing the formation and growth of polygonal ferrite from austenite grain boundaries. The effect is obtained with a content of 0.0003% or more. On the other hand, if the content exceeds 0.0050%, the workability decreases. Therefore, when it contains B, it is set as B: 0.0003% or more and 0.0050% or less of range.

Ca: 0.001% or more and 0.005% or less, REM: One or two kinds selected from 0.001% or more and 0.005% or less Ca and REM spheroidize the shape of the sulfide, and adverse effects of sulfide on stretch flangeability Useful to improve. The effect is obtained when each content is 0.001% or more. On the other hand, if each content exceeds 0.005%, inclusions and the like increase and surface defects and internal defects are caused. Therefore, when Ca and REM are contained, the range is Ca: 0.001% to 0.005% and REM: 0.001% to 0.005%.

In the steel sheet of the present invention, components other than the above are Fe and inevitable impurities. However, as long as the effects of the present invention are not impaired, the inclusion of components other than those described above is not rejected.

Next, the manufacturing method of the high strength steel plate of this invention is demonstrated.
A steel slab adjusted to the above preferred component composition is manufactured, then hot-rolled, and then cold-rolled to obtain a cold-rolled steel sheet. In the present invention, these treatments are not particularly limited, and may be carried out in accordance with conventional methods. Preferred production conditions are as follows. After heating the steel slab to a temperature range of 1000 ° C or higher and 1300 ° C or lower, hot rolling is finished in a temperature range of 870 ° C or higher and 950 ° C or lower, and the obtained hot rolled steel sheet is heated to a temperature of 350 ° C or higher and 720 ° C or lower. Take up in the area. Next, the hot-rolled steel sheet is pickled and then cold-rolled at a rolling reduction in the range of 40% to 90% to obtain a cold-rolled steel sheet.

In the present invention, it is assumed that the steel sheet is manufactured through normal steelmaking, casting, hot rolling, pickling, and cold rolling processes. For example, the steel plate is heated by thin slab casting or strip casting. You may manufacture by omitting a part or all of a hot rolling process.

The obtained cold-rolled steel sheet is subjected to the heat treatment shown in FIG. Hereinafter, a description will be given with reference to FIG.
Annealing is performed for 15 seconds to 1000 seconds in the austenite single phase region. The steel sheet of the present invention is mainly composed of a low-temperature transformation phase obtained by transformation from untransformed austenite, such as martensite, and it is preferable that the polygonal ferrite is as little as possible. is necessary. The annealing temperature is not particularly limited as long as it is in the austenite single phase region, but if the annealing temperature exceeds 1000 ° C., the austenite grain grows remarkably, causing the coarsening of the structure caused by the subsequent cooling and degrading toughness. . Accordingly, the annealing temperature needs to be A ₃ point (austenite transformation point) ° C. or higher, and preferably 1000 ° C. or lower.
Here, A ₃ point is the following formula A ₃ point (° C.) = 910-203 × [C%] ^1/2 + 44.7 × [Si%] − 30 × [Mn%] + 700 × [P%] + 130 x [Al%]
-15.2 × [Ni%]-11 × [Cr%]-20 × [Cu%] + 31.5 × [Mo%] + 104 × [V%]
+ 400 × [Ti%]
Can be calculated. [X%] is mass% of the component element X of the steel sheet.

Also, when the annealing time is less than 15 seconds, the reverse transformation to austenite may not proceed sufficiently, or the carbides in the steel sheet may not be sufficiently dissolved. On the other hand, if the annealing time exceeds 1000 seconds, a cost increase accompanying a great energy consumption is caused. Therefore, the annealing time is in the range of 15 seconds to 1000 seconds. Preferably, it is the range of 60 seconds or more and 500 seconds or less.

The annealed cold-rolled steel sheet is cooled by controlling the average cooling rate to 3 ° C./s or higher to the first temperature range of Ms−150 ° C. or higher and lower than the Ms point. In this cooling, a part of austenite is martensitic transformed by cooling to below the Ms point. Here, if the lower limit of the first temperature range is less than Ms-150 ° C., the amount of untransformed austenite to martensite becomes excessive at this point, and an extremely excellent strength-ductility balance cannot be obtained. On the other hand, when the upper limit of the first temperature range is Ms or more, an appropriate amount of tempered martensite cannot be secured. Therefore, the range of the first temperature range is Ms−150 ° C. or higher and lower than the Ms point. On the other hand, when the average cooling rate is less than 3 ° C./s, excessive formation and growth of polygonal ferrite and precipitation of pearlite occur, and a desired steel sheet structure cannot be obtained. Therefore, the average cooling rate from the annealing temperature to the first temperature range is set to 3 ° C./s or more. Preferably it is 5 degrees C / s or more, More preferably, it is 8 degrees C / s or more. The upper limit of the average cooling rate is not particularly limited as long as the cooling stop temperature does not vary, but in general equipment, when the average cooling rate exceeds 100 ° C / s, the structure in the longitudinal direction and the width direction of the steel plate The dispersion is significantly increased, so that it is preferably 100 ° C./s or less. Therefore, the average cooling rate is preferably in the range of 8 ° C./s to 100 ° C./s.
The Ms point described above is preferably determined by measurement of thermal expansion during cooling by a four-master test or the like, or by actual measurement by measurement of electrical resistance, but can also be obtained by an approximate expression such as the following expression. M is an approximate value obtained empirically.
M point (° C.) = 540−361 × {[C%] / (1− [α%] / 100)} − 6 × [Si%] − 40 × [Mn%]
+ 30 × [Al%] − 20 × [Cr%] − 35 × [V%] − 10 × [Mo%] − 17 × [Ni%]
-10 x [Cu%]
[X%] is the mass% of the component element X of the steel sheet, and [α%] is the area ratio of polygonal ferrite.
The area ratio of polygonal ferrite is measured, for example, by image processing of a 1000 to 3000 times SEM photograph.
Polygonal ferrite is observed in the steel sheet after annealing and cooling under the above-described conditions. For the cold-rolled steel sheet having a desired component composition, the area ratio of polygonal ferrite is determined after annealing and cooling. The value of M can be obtained by substituting into the above equation together with the content of the alloying element obtained from the component composition.

The steel sheet cooled to the first temperature range is heated to a second temperature range of 340 ° C. or more and 520 ° C. or less, and is held for 15 seconds or more and 1000 seconds or less in the second temperature range.
In the second temperature range, martensite generated by cooling from the annealing temperature to the first temperature range is tempered, and the austenite is stabilized by transforming untransformed austenite into upper bainite that suppresses the formation of carbides. When the upper limit of the second temperature range exceeds 520 ° C., a carbide is precipitated from untransformed austenite, so that a desired structure cannot be obtained. On the other hand, when the lower limit of the second temperature range is less than 340 ° C., lower bainite is generated from untransformed austenite, and the amount of C enrichment in the austenite decreases. Therefore, the range of the second temperature range is set to a range of 340 ° C. or more and 520 ° C. or less. Preferably, it is the range of 370 degreeC or more and 450 degrees C or less.

In addition, when the holding time in the second temperature range is less than 15 seconds, the tempering of martensite is insufficient and the desired steel sheet structure cannot be obtained, and as a result, the workability of the obtained steel sheet is sufficiently secured. Therefore, the holding time in the second temperature range needs to be 15 seconds or longer. On the other hand, in the present invention, a holding time in the second temperature range of 1000 seconds is sufficient because of the effect of promoting bainite transformation by martensite generated in the first temperature range even when it is necessary to proceed with the upper bainite transformation. . Usually, when the alloy components such as C, Cr and Mn increase as in the steel of the present invention, the bainite transformation is delayed, but when martensite and untransformed austenite coexist as in the present invention, the bainite transformation rate is remarkably high. There have been some reports on this, and the inventors have also found out about the steel of the present invention. On the other hand, when the holding time in the second temperature range exceeds 1000 seconds, stable residual austenite in which C is concentrated by precipitation of carbides from untransformed austenite which becomes residual austenite as the final structure of the steel sheet cannot be obtained. As a result, the desired strength and / or ductility may not be obtained. Therefore, the holding time is 15 seconds or more and 1000 seconds or less. Preferably, it is 30 seconds or more and 700 seconds or less. More preferably, it is 40 seconds or more and 400 seconds or less.

In the series of heat treatments in the present invention, the holding temperature does not need to be constant as long as it is within the predetermined temperature range described above, and even if it fluctuates within the predetermined temperature range, the gist of the present invention is not impaired. The same applies to the cooling rate. Further, as long as the thermal history is satisfied, the steel sheet may be heat-treated with any equipment. Furthermore, after the heat treatment, it is also included in the scope of the present invention to perform temper rolling on the surface of the steel sheet for shape correction or to perform surface treatment such as electroplating.

In the method for producing a high-strength steel sheet according to the present invention, a hot dip galvanizing treatment or an alloyed hot dip galvanizing treatment obtained by further adding an alloying treatment to the hot dip galvanizing treatment can be added. The hot dip galvanizing treatment or the alloyed hot dip galvanizing treatment may be performed during the temperature rise from the first temperature range to the second temperature range, during the second temperature range hold, or after the second temperature range hold. Even in this case, the holding time in the second temperature range is 15 seconds or more and 1000 seconds or less including the processing time of the hot dip galvanizing treatment or the alloying galvanizing treatment. The hot dip galvanizing treatment or alloying hot dip galvanizing treatment is preferably performed in a continuous hot dip galvanizing line.

Moreover, in the manufacturing method of the high-strength steel sheet of the present invention, after manufacturing the high-strength steel sheet completed up to the heat treatment according to the above-described manufacturing method of the present invention, a hot dip galvanizing process or a further alloying process is performed. Can be added.

The method of performing hot dip galvanizing treatment or alloying hot dip galvanizing treatment on a steel sheet is not particularly limited, and may be performed according to a conventional method. For example, it is as follows.
The steel sheet is infiltrated into the plating bath and the amount of adhesion is adjusted by gas wiping. The amount of dissolved Al in the plating bath should be in the range of 0.12% to 0.22% by mass in the case of hot dip galvanizing, and in the range of 0.08% to 0.18% in the case of galvannealed alloying. preferable.

In the case of hot dip galvanizing, the temperature of the plating bath may be in the range of 450 ° C. to 500 ° C. In the case of further alloying, the temperature during alloying should be 550 ° C. or lower. It is preferable. When the alloying temperature exceeds 550 ° C, carbides precipitate from untransformed austenite, and in some cases pearlite is generated, so strength and workability or both cannot be obtained, and the powdering properties of the plating layer are also low. to degrade. On the other hand, if the temperature during alloying is less than 450 ° C., alloying may not proceed.

The plating adhesion amount is preferably in the range of 20 g / m ² or more and 150 g / m ² or less per side. If the coating amount is less than 20 g / m ² , the corrosion resistance will be insufficient, while if it exceeds 150 g / m ² , the corrosion resistance will be saturated and only increase the cost. The alloying degree of the plating layer (Fe mass% in the plating layer (Fe content)) is preferably in the range of 7 mass% to 15 mass%. If the degree of alloying of the plating layer is less than 7% by mass, unevenness in alloying occurs and the appearance quality deteriorates, or the so-called ζ phase is generated in the plating layer and the slidability of the steel sheet deteriorates. On the other hand, if the degree of alloying of the plating layer exceeds 15% by mass, a large amount of hard and brittle Γ phase is formed, and the plating adhesion deteriorates.

Hereinafter, the present invention will be described in more detail with reference to examples. However, the following examples do not limit the present invention. In addition, changing the configuration within the scope of the gist configuration of the present invention is included in the scope of the present invention.

A steel slab obtained by melting steel having the composition shown in Table 1 is heated to 1200 ° C, and hot-rolled steel sheet that has been hot-rolled and finished at 870 ° C is wound up at 650 ° C, and then the hot-rolled steel sheet is pickled. Thereafter, it was cold-rolled at a rolling rate (rolling rate) of 65% to obtain a cold-rolled steel plate having a thickness of 1.2 mm. The obtained cold-rolled steel sheet was heat-treated under the conditions shown in Table 2. In addition, the cooling stop temperature: T1 in Table 2 is a temperature at which the cooling of the steel sheet is stopped when the steel sheet is cooled from the annealing temperature.

Further, some cold-rolled steel sheets were subjected to hot dip galvanizing treatment or alloying hot dip galvanizing treatment. Here, the hot dip galvanizing treatment was carried out on both sides so that the plating bath temperature was 463 ° C. and the basis weight (per one side) was 50 g / m ² . Also, in the alloying hot dip galvanizing treatment, the plating bath temperature: 463 ° C., the basis weight (per one side): 50 g / m ² , and the degree of alloying (Fe mass% (Fe content)) is 9 mass%. The alloying temperature was adjusted to 550 ° C. or lower and the alloying conditions were adjusted to perform double-sided plating. In addition, the hot dip galvanizing treatment and the alloying hot dip galvanizing treatment were performed after once cooling to T1 ° C. shown in Table 2.

When the obtained steel sheet is not subjected to a plating treatment, the rolling rate (elongation rate) is adjusted to 0.3% after the heat treatment, and when the hot dip galvanizing treatment or the alloyed hot dip galvanizing treatment is performed, after these treatments. Quality rolling was applied.

Various properties of the steel sheet thus obtained were evaluated by the following methods.
Samples are cut out from each steel plate, polished, and the surface having a normal parallel to the plate width direction is observed with a scanning electron microscope (SEM) at 10 000 magnifications to measure the area ratio of each phase. The phase structure of each crystal grain was identified.

The method for measuring the total length of the prior austenite grain boundaries and the method of obtaining the ratio of the prior austenite grain boundaries in the tempered martensite or adjacent to the tempered martensite are as follows.
Former austenite grain boundaries are etched with a corrosive solution prepared by mixing HCl and glycerin as a reaction rate adjusting agent in picric acid + surfactant + ferrous chloride + oxalic acid disclosed in JP-A-2005-241635. Made it appear. From this, the old austenite grain boundaries were observed with an optical microscope at × 500 to × 1000 times, the total length was measured with an image processing device, and then the structure of the same field of view that was mirror polished and nital etched was again observed with an SEM. The ratio of the prior austenite grain boundaries inside or adjacent to the tempered martensite was determined.

The amount of retained austenite was determined by measuring the X-ray diffraction intensity after grinding and polishing the steel plate to 1/4 of the plate thickness in the plate thickness direction. Co-Kα is used for incident X-rays, and the residual from the intensity ratio of each surface of austenite (200), (220), (311) to the diffraction intensity of each surface of ferrite (200), (211), (220) The amount of austenite was calculated.

The average amount of C in the retained austenite is obtained by calculating the lattice constant from the intensity peaks of the (200), (220) and (311) surfaces of austenite in the X-ray diffraction intensity measurement. C (mass%) was determined.
[C%] = (a ₀ -0.3580-0.00095 × [Mn%]-0.0056 × [Al%]-0.022 × [N%]) / 0.0033
However, a _0: the lattice constant (nm), [X%] : % by weight of the element X. In addition, mass% of elements other than C was mass% with respect to the whole steel plate.

Tensile test was conducted using JIS No. 5 test piece (JIS
Z 2201) was used in accordance with JIS Z 2241. TS (tensile strength) and T.EL (total elongation) were measured, and the product of strength and total elongation (TS × T.EL) was calculated to evaluate the balance between strength and workability (ductility). In the present invention, the case of TS × T.EL ≧ 29000 (MPa ·%) was considered good. The above evaluation results are shown in Tables 3-1 and 3-2.

As is apparent from Table 3, all the steel sheets of the present invention have a high strength and excellent workability with a tensile strength of 1470 MPa or more and a TS × T.EL value of 29000 MPa ·% or more, It was confirmed that the ductility was excellent.

In accordance with the present invention, after the addition of Cr having a C content in the steel sheet of 0.30% or more and a ferrite suppressing effect, the steel sheet annealed in the austenite single phase region is rapidly cooled to partially martensite the austenite. By tempering martensite and stabilizing retained austenite, it is possible to obtain a high-strength steel sheet that is remarkably excellent in workability, particularly the balance between strength and ductility, and that has a tensile strength of 1470 MPa or more.

Claims (11)

In mass% C: 0.30% or more and 0.73% or less,
Si: 3.0% or less,
Al: 3.0% or less,
Si + Al: 0.7% or more,
Cr: 0.2% to 8.0%,
Mn: 10.0% or less,
Cr + Mn: 1.0% or more
P: 0.1% or less,
S: 0.07% or less and N: 0.010% or less, the balance is composed of Fe and inevitable impurities,
As the steel sheet structure, the area ratio of martensite to the entire steel sheet structure is 15% to 90%, the amount of retained austenite is 10% to 50%, and 50% or more of the martensite is tempered martensite and the tempered martensite. Satisfies the area ratio of the site steel sheet structure to 10% or more, the area ratio of polygonal ferrite to the entire steel sheet structure of 10% or less (including 0%), tensile strength of 1470 MPa or more, tensile strength x total elongation Is a high-strength steel sheet characterized by a 29000 MPa ·% or more. The high-strength steel sheet according to claim 1, wherein 30% or more of the total length of the prior austenite grain boundary is present in the tempered martensite or adjacent to the tempered martensite. The high-strength steel sheet according to claim 1 or 2, wherein an average amount of C in the retained austenite is 0.7 mass% or more. The steel sheet is further in mass%,
The high strength according to any one of claims 1 to 3, characterized by containing Ni: 0.05% or more and 5.0% or less, and satisfying Cr + Mn + Ni: 1.0% or more instead of Cr + Mn: 1.0% or more. steel sheet. The steel sheet is further in mass%,
V: 0.005% to 1.0%,
The composition according to any one of claims 1 to 4, characterized by containing one or more selected from Mo: 0.005% to 0.5% and Cu: 0.05% to 2.0%. High strength steel plate. The steel sheet is further in mass%,
The high content according to any one of claims 1 to 5, characterized by containing one or two selected from Ti: 0.01% to 0.1% and Nb: 0.01% to 0.1%. Strength steel plate. The steel sheet is further in mass%,
The high-strength steel sheet according to any one of claims 1 to 6, wherein B: 0.0003% or more and 0.0050% or less is contained. The steel sheet is further in mass%,
The high content according to any one of claims 1 to 7, comprising one or two selected from Ca: 0.001% to 0.005% and REM: 0.001% to 0.005%. Strength steel plate. A high-strength steel sheet comprising a hot-dip galvanized layer or an alloyed hot-dip galvanized layer on the surface of the steel sheet according to any one of claims 1 to 8. A steel slab having the composition according to any one of claims 1 to 8 is hot-rolled and then cold-rolled into a cold-rolled steel sheet, and then the cold-rolled steel sheet is at least 15 seconds in the austenite single-phase region. After annealing for 1000 seconds or less, cool at an average cooling rate of 3 ° C / s or higher up to the first temperature range of Ms-150 ° C or higher and lower than Ms with respect to the martensite transformation start temperature Ms, and then 340 ° C or higher and 520 ° C. A method for producing a high-strength steel sheet, characterized in that the temperature is raised to the following second temperature range, and subsequently maintained in the second temperature range for 15 seconds to 1000 seconds. 11. The high-strength steel sheet according to claim 10, wherein hot-dip galvanizing treatment or alloying hot-dip galvanizing treatment is performed during temperature rising to the second temperature range or holding in the second temperature range. Manufacturing method.

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